- Email: [email protected]
The effects of subsurface microstructure evolution on fretting wear resistance of nickel-based alloy
Wear 416–417 (2018) 81–88
Contents lists available at ScienceDirect
Wear journal homepage: www.elsevier.com/locate/wear
The effects of subsurface microstructure evolution on fretting wear resistance of nickel-based alloy
T
⁎
J. Lia, Y.H. Lub, , X.H. Tua, W. Lia a b
Institute of Advanced Wear & Corrosion Resistant and Functional Materials, Jinan University, Guangzhou 510632, China National Center for Materials Service Safety, University of Science and Technology Beijing, Beijing 100083, China
A R T I C LE I N FO
A B S T R A C T
Keywords: Non-ferrous metals Nanostructure Wear
In order to clarify the reason for the differences in the fretting wear resistance of Inconel 690 and 600 plate against 304 stainless steel ball, which is conducive to deepening the understanding of degradation mechanism, the comparative study on wear scars of two materials have been conducted in the atmosphere. The results indicated that Inconel 690 exhibited more excellent wear resistance than Inconel 600, which could be attributed to the thicker and harder tribological transformed structure (TTS) formed in worn subsurface of Inconel 690. The further structural characterizations of subsurface showed that the transformation mechanism from the matrix (coarse grains) to TTS (nano-sized grains) was dynamic recrystallization. Compared with Inconel 600, Inconel 690 possessed lower stacking fault energy (SFE) for its high chromium content, which prompted the occurrence of dynamic recrystallization and eventually led to a harder and thicker TTS in Inconel 690 worn subsurface.
1. Introduction Nickel-based alloys 690 and 600 (Inconel 690, 600) are extensively used for steam generator (SG) tube material due to their excellent corrosion resistance. Up to now, the failure researches on SG tube have mainly focused on the corrosion related problems. However, fretting wear, which occurs at the contact surface of SG tube and its support plate due to flow-induced vibrations, has become one of the most important problems in pressurized water reactor nuclear power plants with the increasing of service life [1,2]. The surface damage induced by fretting wear could cause the thinning and rupture of SG tube, and eventually threaten the safety of nuclear power plants. The agreement has been reached that although the experimental environment differs greatly from the working environment of SG tube, the fundamental work could play a positive role in deepening the understanding of degradation mechanism during fretting wear. Therefore, many fundamental investigations have been conducted in the air or pure water to research the fretting wear behavior of SG tube. Among them, the early basic research works have concentrated on revealing the wear laws related to mechanical factors (such as displacement amplitude, normal force, frequency) as they vary with the contact position due to the fluctuation of flow velocity in SG tube, and establishing the wear mechanism through the characterization of the worn scar surface [3–9]. For example, the wear volume decreases with the
⁎
increasing of normal force as the wear mechanisms change from oxidation wear to adhesive wear [7]. The comparative studies on the fretting wear behavior between Inconel 600 and 690 have also been reported. However, the underlying causes for the difference in fretting wear resistance between the two materials are far from being clearly revealed. Recently, more and more researches indicate that the material's fretting wear resistance is closely related to the subsurface deformation behavior [10–14]. For example, Inconel 690 with small grains presents more excellent fretting wear resistance than that with larger ones. The reason for this is that the reduction in grain size favors the formation of tribological transformed structure (TTS) in the worn subsurface. TTS could improve the wear resistance of materials due to its ultra-high hardness [11]. The reason why the wear volume of Inconel 690 induced by fretting wear in cathodic condition is smaller than that in anodic condition is that the oxidation of TTS is suppressed in cathodic condition [14]. With the aim to improve the understanding of degradation mechanism during fretting wear, comprehensive characterizations of worn surface and subsurface are firstly conducted and analyzed. Then a schematic diagram for the formation of TTS in worn scar subsurface is proposed. Finally, the effects of subsurface deformation on fretting wear behavior of Inconel 690 and 600 are discussed.
Corresponding author. E-mail address: [email protected] (Y.H. Lu).
https://doi.org/10.1016/j.wear.2018.10.003 Received 12 June 2018; Received in revised form 28 August 2018; Accepted 1 October 2018 Available online 11 October 2018 0043-1648/ © 2018 Published by Elsevier B.V.
Wear 416–417 (2018) 81–88
J. Li et al.
Table 1 Chemical compositions of Inconel 600 and 690. Specimen
Element (wt%)
Inconel 600 Inconel 690 304 SS
Ni
Fe
Cr
C
Ti
Mn
Si
N
S
P
Mo
Bal Bal 9.35
11.0 11.6 Bal
15.2 29.9 18.3
0.022 0.025 0.018
0.30 0.30 –
0.23 0.25 0.25
0.29 0.33 0.31
0.024 0.020 –
0.025 0.025 0.025
0.086 0.01 0.034
– – –
Fig. 1. The microstructure of (a) Inconel 600 and (b) Inconel 690.
Fig. 2. (a) The coefficient of friction and (b) the typical cross-sectional profiles of worn scars in the case of Inconel 600 and 690.
2. Experiment
displacement amplitude at 100 µm, frequency at 20 Hz, test duration at 30 min. By calculation, the initial maximal Hertz contact pressure was about 2.18 GPa. The chosen values of test parameter had referred to the previous fundamental investigation [3,5]. Each test was repeated at least two times to check the repeatability. All specimens were ground and polished to get mirror finish before test. After fretting wear test, the loose debris was removed by compressed air and ultrasonic cleaning in alcohol for 10 min. The wear volume and profiles of worn scars were firstly measured by a 3D surface profiler (America ADE MicroXAM-3D). The worn surfaces and subsurface morphologies were observed using the field emission scanning electron microscope (FESEM) equipped with energy dispersive X-ray (EDX) detector for element analysis. Then, the nanoindentation tests were performed on microstructures in worn subsurface. On each specimen, hardness tests were carried out three times at the same depth. Finally, the TEM lamellas were prepared by ZEISS Auriga focused ion beam (FIB) equipment through the procedure of depositing, milling and polishing. The worn scar subsurface microstructures of Inconel 690 were characterized by transmission electron
The friction pair for this study was 304 stainless steel (304 SS) balls and Inconel 690 or 600 plates. The chemical compositions of the tribomaterials were listed in Table 1. It could be seen that the nickel-based Inconel 690 and 600 were mainly composed of nickel, chromium and iron element. The optical microstructures of Inconel 600 and 690 were shown in Fig. 1a and b, which had been revealed by electrolytic etching in 10% oxalate solution at 5 V for 25 s. The average grain size of Inconel 600 calculated by 200 grains was about 17.4 µm with bulk hardness of 240 HV. Correspondingly, they are 28 µm and 200 HV for Inconel 690. Fretting wear tests were performed on an Optimol SRV IV oscillating friction tester using a ball-on-flat contact configuration at room temperature (25–30 °C) in air condition. The upper specimen was 304SS ball with a diameter of 10 mm, and the lower specimen was cut to a dimension of 15 mm × 15 mm × 2.1 mm then conglutinated to a stainless base. Details of this device had been described in the previous study [11]. The test parameters were set as normal load at 60 N, 82
Wear 416–417 (2018) 81–88
J. Li et al.
Fig. 3. The SEM morphologies as well as oxygen distribution of worn scar surface: (a) Inconel 600 and (b) Inconel 690.
deformed layer (GDL, between dash dot- long dash dot lines). The grain boundaries are the criteria for differentiating TTS and GDL as they cannot be observed in TTS but are deformed and parallel to the fretting direction in GDL. However, it should be mentioned that the WDL thickness of Inconel 600 is about 5 µm, which is larger than that of Inconel 690 (~ 3 µm). But the TTS thickness of Inconel 600 (~ 10 µm) is much thinner than that of Inconel 690 (~ 20 µm). Finally, there is no noticeable difference in GDL thickness between Inconel 600 and 690. Fig. 5 shows the load-indentation depth curve of different microstructure in worn subsurface. The locations tested were about 6 µm (TTS), 26 µm (GDL) and 46 µm (Matrix) from the surface, respectively. The hardness of TTS, GDL and Matrix is 5.33, 4.54 and 3.89 GPa in worn subsurface of Inconel 690. Correspondingly, they are 4.95, and 4.43 and 4.09 GPa in the case of Inconel 600. It is obvious that the hardness decreases from TTS to GDL to Matrix in both cases. In addition, the hardness of TTS and GDL in worn subsurface of Inconel 690 is slightly larger than that of Inconel 600. In order to find out the reason why the thickness and hardness of stratified structures varies with Inconel 690 and 600, the microstructures of GDL and TTS in nanoscale are characterized. Firstly, two TEM samples are obtained from Inconel 690 worn subsurface using FIB. As shown in Fig. 6, sample 1 is taken from the interface of GDL/Matrix, using an original grain boundary (GB) as the reference. Therefore, the section below the original GB is Matrix, the upper is GDL. And sample 2 is taken from the interface of TTS/GDL, using a deformed GB as the reference. Correspondingly, the section below the deformed GB is GDL, the upper section is TTS. To facilitate the following analysis, the directions from Matrix to GDL as well as GDL to TTS are expressed using a to b and b to c, respectively. Fig. 7 shows the characterizations of sample 1 located at the interface of GDL/Matrix. From the bright field TEM (BFTEM) image shown in Fig. 7a, the material above the dash line is platinum, which is deposited and used as the protective layer during the preparation of TEM sample. As a result, the sample could be divided into two parts according to the GB: Grain 1 and Grain 2. Combined with the selected
microscope (TEM) using FEI Tecnai F20. 3. Results Fig. 2a shows the coefficient of friction (COF) as a function of cycle number during fretting wear in the case of Inconel 600 and 690. It could be found that the COF of both cases increases rapidly at the initial stage, and then gradually tends to a steady state with fluctuations. However, although the average COF of both is about 1, the COF of Inconel 600 presents a larger fluctuation than that of Inconel 690. Fig. 2b shows the typical cross-sectional profiles of worn scar along the fretting direction. It could be seen that a U-shape worn scar is produced in both cases. By comparison, the worn scar of Inconel 600 is deeper and wider than that of Inconel 690. Correspondingly, the wear volume is 58.5 × 105 μm3 for Inconel 600 and 47.5 × 105 μm3 for Inconel 690. In fact, the result of parallel experiments indicates that the average wear volume of Inconel 600 is about 1.15 times larger than that of Inconel 690. Fig. 3a and b show the worn scar surface SEM images and the oxygen distribution (the upper right corner) of Inconel 600 and 690, respectively. According to the characterization shown in Fig. 3a and b, in both cases, the oxygen content inside the worn scar is larger than that outside the worn scar, which indicates that the severe oxidation behaviors have occurred on the whole worn scar surface. Then, the enlarged SEM morphologies of position A and B show that both the worn scar surfaces are covered by continuous wear debris. As a result, there is no obvious difference on worn scar surface between Inconel 690 and 600. For further research, the etched worn subsurface SEM images of Inconel 600 and 690 are characterized and shown in Fig. 4a and b, respectively. The locations of worn subsurface shown in Fig. 4a and b correspond to the marks A and B in Fig. 3. It could be seen that there is a stratified structures in the worn subsurface. From the outmost surface to material inside, the structures are wear debris layer (WDL, between dash-dash lines), TTS (between dash-dash dot lines) and general 83
Wear 416–417 (2018) 81–88
J. Li et al.
Fig. 4. The SEM morphologies worn scar subsurface: (a) Inconel 600 and (b) Inconel 690.
brightness of the sample is extremely uneven. According to the block's shape, the sample can be divided into two parts. The left part of dash dot line is equiaxed blocks and the right part is striplike blocks. From the BFTEM image of C position shown in Fig. 8b, it could be found that the striplike blocks correspond to some fine twins and subgrains. They are dozens of nanometers in width and hundreds of nanometers in length. Meanwhile, some dislocation arrays and walls are observed inside the twins and subgrains. As shown in Fig. 8c, the high resolution transmission electron microscope (HRTEM) image of C1 indicates the observed twins could be the mechanical twins due to the dislocations and steps located at the twin boundary, which are the unique characteristics of mechanical twins compared with annealing or growth twins [15]. According to the BFTEM and concentric SADP images of D position shown in Fig. 8d, the equiaxed blocks correspond to randomly oriented equiaxed nano-sized grains. The statistic distribution of equiaxed grain sizes in Fig. 8e shows that the grain size ranges from 20 nm to 90 nm with the average of 47.3 nm.
Fig. 5. The load-indentation depth curve of different microstructure in worn subsurface of Inconel 690 and 600.
4. Discussion In order to increase the thermal conductivity, the thickness of SG tube was reduced from 2 mm to 1.09 mm. As a result, the wear depth was also an importance factor for reliability of SG tube during fretting wear test. As shown in Fig. 2b, both the deeper worn scar and the larger wear volume indicated that Inconel 690 exhibited more excellent wear resistance than Inconel 600. By comparing the morphologies of wear scar surface and subsurface, there was no significant difference on worn scar surfaces as they were both covered by the continuous wear debris as shown in Fig. 3. From the morphologies shown in Fig. 4, fretting wear caused the formation of stratified structures in the worn subsurface, which were WDL, TTS, GDL and matrix from outside into inside, respectively. And it was obvious that Inconel 690 possessed a thinner WDL but a thicker TTS compared with Inconel 600. The thinner WDL of Inconel 690 could be attributed to the lower wear volume for its excellent wear resistance. According to other's previous studies, the thicker dense and continuous WDL could be more effective to separate the contacting surface and act as a lubricant. To some extent, it has benefits for the improvement of material's wear resistance [16–18]. However, compared with Inconel 600, Inconel 690 with thinner WDL but excellent wear resistance indicated that the wear resistance was not primarily dependent on WDL in this research. Therefore, the reason why Inconel 690 presented more excellent fretting wear resistance than Inconel 600 could depends mainly on TTS as it was directly related to material loss, which could be explained as follows. On one hand, it had been reported that the wear resistance was proportional to the hardness of material within a certain range [19]. The TTS's hardness of Inconel 690 was greater than that of Inconel 600 as shown in Fig. 5, resulting in an increase in wear resistance. On the other hand, the increasing in the thickness of the deformation layer would reduce the transfer of stress
Fig. 6. The schematic view of FIB sampling position at Inconel 690 worn subsurface.
area diffraction pattern (SADP) of A and B as shown in Figs. 7b and 7c, the high-density tangled dislocations are formed in Grain 1. Meanwhile, the multiple twins associated with high-density tangled dislocations are formed in the Grain 2. The length of twins is up to several micrometers, but the width of twins is only hundreds of nanometers. From the dark field TEM (DFTEM) of A and B shown in Figs. 7d and 7e, the brightness in Grain 1 or Grain 2 are inconsistent, which indicates that the grains have been subdivided into several regions with different crystallographic orientations. Fig. 8 shows the characterizations of sample 2 located at the interface of TTS/GDL. From the DFTEM image shown in Fig. 8a, the
84
Wear 416–417 (2018) 81–88
J. Li et al.
Fig. 7. The characterization of sample 1 located at the interface of GDL/Matrix: (a) overall BFTEM image, (b) SADP of A position, (c) SADP of B position, (d) DFTEM of A position and (e) DFTEM of B position.
85
Wear 416–417 (2018) 81–88
J. Li et al.
Fig. 8. The characterizations of sample 2 located at the interface of TTS/GDL: (a) the overall DFTEM image, (b) the BFTEM of C position, (c) the HRTEM of C1 position, (d) the BFTEM as well as SADP images of D position and (e) the statistics of grain size in TTS.
gradually from contact surface to material inside. In order to accommodate plastic strains in polycrystalline materials, various dislocation activities were normally motivated in low strain conditions [21,22]. As a result, there were plenty of tangled dislocations in matrix (I stage, Fig. 7-Grain 1). Secondly, twinning was an important way of plastic deformation, which often acted as a supplement when dislocation slips was not easy to occur [23,24]. In addition, experimental results indicate that the critical stress for twinning was larger than dislocation slip [25]. Thence, many mechanical twins associated with dislocations were formed in coarse grain (II stage, Fig. 7-Grain 2). Furthermore, the twins would continue to undergo twin deformation with the further increase of strain, resulting in a reduction of the thickness and length of twins
from the surface to the substrate [20]. This would also enhance the material's wear resistance. As mentioned above, the wear resistance was closely related to the thickness and hardness of TTS. As a result, the relationship of TTS's thickness and hardness with nickel-based alloy would be discussed based on the formation mechanism of TTS. The first was about the formation mechanism of TTS. According to the microstructure characterization shown in Figs. 7 and 8, the transformation from matrix to TTS could be proposed using the schematic diagram shown in Fig. 9. According to previous research results of others, the formation mechanism of TTS could be analyzed as follows. Firstly, it is well known that the strain could be generated during fretting wear and decrease 86
Wear 416–417 (2018) 81–88
J. Li et al.
Fig. 9. Schematic diagram for TTS formation.
COF on worn scar surface of Inconel 600 was greater than that of Inconel 690.
and subgrains (III stage, Fig. 8b). In addition, including the strain, friction heat was also generated on contact surface during fretting wear. According to the recrystallization theory, the increase of temperature could cause the formation of dislocation walls as shown in Fig. 8b. With the further increase of strain and temperature, the tangled dislocation and dislocation walls could act as nuclei. Finally, the striplike twins and subgrains transformed into equiaxed grain (IV stage, Fig. 8d) as friction heat and strain increased and reached to the threshold value for dynamic recrystallization (DC). In summary, in nanoscale, GDL was mainly composed of striplike twins and subgrains, while TTS consisted of equiaxed nanosized grains. The transformation mechanism from the matrix (coarse grains) to TTS (nano-sized grains) was DC. The formation of TTS was in the form of DC according to the above discussion. Then, the relationship between DC and nickel-based alloy should be revealed. It is well known that the DC is closely related to material's stacking fault energy (SFE). Materials with lower SFE were usually prone to DC as the dislocation climb and cross slip were difficult to occur [26–28]. In addition, the research about the effect of chromium content on SFE showed that SFE rapidly decreased with the increase of chromium content in nickel-base alloys [29]. From the chemical composition shown in Table 1, the chromium content of Inconel 690 was about 30 wt%, which was as twice as that of Inconel 600. In short, the higher chromium content led to the lower SFE, which prompted the occurrence of DC and eventually created a thicker and harder TTS in Inconel 690 worn subsurface. As a result, compared with Inconel 600, Inconel 690 with thicker TTS exhibited more excellent wear resistance. Furthermore, it should be mentioned that the reduction of grain size was conducive to the formation of TTS due to the increasing of bulk hardness, which had been proposed in our previous work [11]. However, the bulk hardness of Inconel 690 was smaller than that of Inconel 600, but had more excellent wear resistance in present study. As a result, compared with grain size, the Cr content could present a greater impact on the thickness of TTS. As for the COF shown in Fig. 2a, it was reported that the fluctuation of COF was closely related to the formation of wear debris and their ejection from the contact area [30]. Compared with Inconel 690, more wear debris generated during fretting wear of Inconel 600, which could be deduced from the larger wear volume. Therefore, the fluctuation of
5. Conclusion The fretting wear behaviors of Inconel 600 and 690 in the atmosphere were investigated using the 3D surface profiler, SEM, EDX and TEM. The observation and analysis led to the following conclusions: (1) Fretting wear caused the formation of stratified structures in the worn subsurface of Inconel 690 and 600, which were WDL, TTS, GDL and matrix from outside into inside, respectively. (2) Compared with Inconel 600, a thinner WDL associated with a thicker and harder TTS were formed in worn subsurface of Inconel 690. (3) The transformation mechanism from the matrix (coarse grains) to TTS (nano-sized grains) was dynamic recrystallization. (4) Inconel 690 exhibited more excellent wear resistance than Inconel 600, which was mainly attributed to the thicker and harder TTS in the worn subsurface of Inconel 690 for its high chromium content. Acknowledgements This work was supported by the National Key Basic Research Program of China (973 Program) (No. 2011CB610504) and the Guangdong Province Science and Technology Plan (No. 2017B090903005). References [1] M.H. Attia, Fretting fatigue and wear damage of structural components in nuclear power stations-fitness for service and life management perspective, Tribol. Int. 39 (2006) 1294–1304. [2] F.M. Guérout, N.J. Fisher, Steam generator fretting-wear damage: a summary of recent findings, J. Press. Vessel. Technol. 121 (1999) 304–310. [3] M.K. Lim, S.D. Oh, Y.Z. Lee, Friction and wear of Inconel 690TT and Inconel 600 for steam generator tube in room temperature water, Nucl. Eng. Des. 226 (2003) 97–105. [4] Y. Yoon, I. Etsion, F.E. Talke, The evolution of fretting wear in a micro-spherical contact, Wear 270 (2011) 567–575.
87
Wear 416–417 (2018) 81–88
J. Li et al.
steel contacts, Tribol. Int. 44 (2011) 1452–1460. [19] Y.S. Zhang, Z. Han, K. Lu, Fretting wear behavior of nanocrystalline surface layer of copper under dry condition, Wear 265 (2008) 396–401. [20] S. Fayeulle, P. Blanchard, L. Vincent, Fretting behavior of titanium alloys, Tribol. T. 36 (1993) 267–275. [21] B.J. Griffiths, D.C. Furze, Tribological Advantages of White Layers Produced by Machining, J. Tribol. 109 (1987) 338–342. [22] W.L. Chan, M.W. Fu, J. Lu, The size effect on micro deformation behavior in microscale plastic deformation, Mater. Des. 32 (2011) 198–206. [23] Z. Wu, H. Bei, F. Otto, G.M. Pharr, E.P. George, Recovery, recrystallization, grain growth and phase stability of a family of FCC-structured multi-component equiatomic solid solution alloys, Intermetallics 46 (2014) 131–140. [24] C. Zheng, D. Raabe, Interaction between recrystallization and phase transformation during intercritical annealing in a cold-rolled dual-phase steel: a cellular automaton model, Acta Mater. 61 (2013) 5504–5517. [25] Y.T. Zhu, X.Z. Liao, X.L. Wu, J. Narayan, Grain size effect on deformation twinning and detwinning, J. Mater. Sci. 48 (2013) 4467–4475. [26] S.L. Guo, D.F. Li, Z.G. Wu, Constitutive relationship and critical condition for dynamic recrystallization of inconel 625 during hot deformation, Adv. Mater. Res. 538–541 (2012) 1240–1244. [27] M.S. Chen, Y.C. Lin, X.S. Ma, The kinetics of dynamic recrystallization of 42CrMo steel, Mater. Sci. Eng. A 556 (2012) 260–266. [28] N.R. Tao, X.L. Wu, M.L. Sui, Grain refinement at the nanoscale via mechanical twinning and dislocation interaction in a nickel-based alloy, J. Mater. Res. 19 (2004) 1623–1629. [29] D.M. Symons, Hydrogen embrittlement of Ni-Cr-Fe alloys, Metall. Mater. Trans. A 28 (1997) 655–663. [30] N. Diomidis, S. Mischler, Third body effects on friction and wear during fretting of steel contacts, Tribol. Int. 44 (2011) 1452–1460.
[5] J.Y. Yun, M.C. Park, G.S. Shin, Effects of amplitude and frequency on the wear mode change of Inconel 690TT SG tube mated with SUS 409, Wear 313 (2014) 83–88. [6] Z.H. Wang, Y.H. Lu, J. Li, T. Shoji, Effect of pH value on the fretting wear behavior of Inconel 690 alloy, Tribol. Int. 95 (2016) 162–169. [7] J. Li, Y.H. Lu, Effects of displacement amplitude on fretting wear behaviors and mechanism of Inconel 600 alloy, Wear 304 (2013) 223–230. [8] X.Y. Zhang, P.D. Ren, J.F. Peng, Fretting wear behavior of Inconel 690 in hydrazine environments, Trans. Nonferrous Met. Soc. China 24 (2014) 360–367. [9] Y. Lee, H. Kim, A comparative study on the fretting wear of steam generator tubes in Korean power plants, Wear 255 (2003) 1198–1208. [10] L. Xin, B.B. Yang, J. Li, Z.H. Wang, Y.H. Lu, T. Shoji, Microstructural evolution of subsurface on Inconel 690TT alloy subjected to fretting wear at elevated temperature, Mater. Des. 104 (2016) 152–161. [11] J. Li, Y.H. Lu, H.Y. Zhang, L. Xin, Effect of grain size and hardness on fretting wear behavior of Inconel 600 alloys, Tribol. Int. 81 (2015) 215–222. [12] M.A.L. Tobi, J. Ding, G. Bandak, S.B. Leen, P.H. Shipway, A study on the interaction between fretting wear and cyclic plasticity for Ti-6Al-4V, Wear 267 (2009) 270–282. [13] B. Yao, Z. Han, K. Lu, Correlation between wear resistance and subsurface recrystallization structure in copper, Wear 3 (2012) 438–445. [14] J. Li, B.B. Yang, Y.H. Lu, L. Xin, Z.H. Wang, T. Shoji, The effects of electrochemical polarization condition and applied potential on tribocorrosion behaviors of Inconel 690 alloys in water environment, Mater. Des. 119 (2017) 93–103. [15] L. Lu, Y. Shen, X. Chen, L. Qian, K. Lu, Ultrahigh strength and high electrical conductivity in copper, Science 304 (2004) 422–426. [16] M. Verenberg, G. Halperin, I. Etsion, Different aspects of the role of wear debris in frettingwear, Wear 252 (2002) 902–910. [17] J. Warburton, The fretting of mild steel in air, Wear 131 (1989) 365–386. [18] N. Diomidis, S. Mischler, Third body effects on friction and wear during fretting of
88
Contents lists available at ScienceDirect
Wear journal homepage: www.elsevier.com/locate/wear
The effects of subsurface microstructure evolution on fretting wear resistance of nickel-based alloy
T
⁎
J. Lia, Y.H. Lub, , X.H. Tua, W. Lia a b
Institute of Advanced Wear & Corrosion Resistant and Functional Materials, Jinan University, Guangzhou 510632, China National Center for Materials Service Safety, University of Science and Technology Beijing, Beijing 100083, China
A R T I C LE I N FO
A B S T R A C T
Keywords: Non-ferrous metals Nanostructure Wear
In order to clarify the reason for the differences in the fretting wear resistance of Inconel 690 and 600 plate against 304 stainless steel ball, which is conducive to deepening the understanding of degradation mechanism, the comparative study on wear scars of two materials have been conducted in the atmosphere. The results indicated that Inconel 690 exhibited more excellent wear resistance than Inconel 600, which could be attributed to the thicker and harder tribological transformed structure (TTS) formed in worn subsurface of Inconel 690. The further structural characterizations of subsurface showed that the transformation mechanism from the matrix (coarse grains) to TTS (nano-sized grains) was dynamic recrystallization. Compared with Inconel 600, Inconel 690 possessed lower stacking fault energy (SFE) for its high chromium content, which prompted the occurrence of dynamic recrystallization and eventually led to a harder and thicker TTS in Inconel 690 worn subsurface.
1. Introduction Nickel-based alloys 690 and 600 (Inconel 690, 600) are extensively used for steam generator (SG) tube material due to their excellent corrosion resistance. Up to now, the failure researches on SG tube have mainly focused on the corrosion related problems. However, fretting wear, which occurs at the contact surface of SG tube and its support plate due to flow-induced vibrations, has become one of the most important problems in pressurized water reactor nuclear power plants with the increasing of service life [1,2]. The surface damage induced by fretting wear could cause the thinning and rupture of SG tube, and eventually threaten the safety of nuclear power plants. The agreement has been reached that although the experimental environment differs greatly from the working environment of SG tube, the fundamental work could play a positive role in deepening the understanding of degradation mechanism during fretting wear. Therefore, many fundamental investigations have been conducted in the air or pure water to research the fretting wear behavior of SG tube. Among them, the early basic research works have concentrated on revealing the wear laws related to mechanical factors (such as displacement amplitude, normal force, frequency) as they vary with the contact position due to the fluctuation of flow velocity in SG tube, and establishing the wear mechanism through the characterization of the worn scar surface [3–9]. For example, the wear volume decreases with the
⁎
increasing of normal force as the wear mechanisms change from oxidation wear to adhesive wear [7]. The comparative studies on the fretting wear behavior between Inconel 600 and 690 have also been reported. However, the underlying causes for the difference in fretting wear resistance between the two materials are far from being clearly revealed. Recently, more and more researches indicate that the material's fretting wear resistance is closely related to the subsurface deformation behavior [10–14]. For example, Inconel 690 with small grains presents more excellent fretting wear resistance than that with larger ones. The reason for this is that the reduction in grain size favors the formation of tribological transformed structure (TTS) in the worn subsurface. TTS could improve the wear resistance of materials due to its ultra-high hardness [11]. The reason why the wear volume of Inconel 690 induced by fretting wear in cathodic condition is smaller than that in anodic condition is that the oxidation of TTS is suppressed in cathodic condition [14]. With the aim to improve the understanding of degradation mechanism during fretting wear, comprehensive characterizations of worn surface and subsurface are firstly conducted and analyzed. Then a schematic diagram for the formation of TTS in worn scar subsurface is proposed. Finally, the effects of subsurface deformation on fretting wear behavior of Inconel 690 and 600 are discussed.
Corresponding author. E-mail address: [email protected] (Y.H. Lu).
https://doi.org/10.1016/j.wear.2018.10.003 Received 12 June 2018; Received in revised form 28 August 2018; Accepted 1 October 2018 Available online 11 October 2018 0043-1648/ © 2018 Published by Elsevier B.V.
Wear 416–417 (2018) 81–88
J. Li et al.
Table 1 Chemical compositions of Inconel 600 and 690. Specimen
Element (wt%)
Inconel 600 Inconel 690 304 SS
Ni
Fe
Cr
C
Ti
Mn
Si
N
S
P
Mo
Bal Bal 9.35
11.0 11.6 Bal
15.2 29.9 18.3
0.022 0.025 0.018
0.30 0.30 –
0.23 0.25 0.25
0.29 0.33 0.31
0.024 0.020 –
0.025 0.025 0.025
0.086 0.01 0.034
– – –
Fig. 1. The microstructure of (a) Inconel 600 and (b) Inconel 690.
Fig. 2. (a) The coefficient of friction and (b) the typical cross-sectional profiles of worn scars in the case of Inconel 600 and 690.
2. Experiment
displacement amplitude at 100 µm, frequency at 20 Hz, test duration at 30 min. By calculation, the initial maximal Hertz contact pressure was about 2.18 GPa. The chosen values of test parameter had referred to the previous fundamental investigation [3,5]. Each test was repeated at least two times to check the repeatability. All specimens were ground and polished to get mirror finish before test. After fretting wear test, the loose debris was removed by compressed air and ultrasonic cleaning in alcohol for 10 min. The wear volume and profiles of worn scars were firstly measured by a 3D surface profiler (America ADE MicroXAM-3D). The worn surfaces and subsurface morphologies were observed using the field emission scanning electron microscope (FESEM) equipped with energy dispersive X-ray (EDX) detector for element analysis. Then, the nanoindentation tests were performed on microstructures in worn subsurface. On each specimen, hardness tests were carried out three times at the same depth. Finally, the TEM lamellas were prepared by ZEISS Auriga focused ion beam (FIB) equipment through the procedure of depositing, milling and polishing. The worn scar subsurface microstructures of Inconel 690 were characterized by transmission electron
The friction pair for this study was 304 stainless steel (304 SS) balls and Inconel 690 or 600 plates. The chemical compositions of the tribomaterials were listed in Table 1. It could be seen that the nickel-based Inconel 690 and 600 were mainly composed of nickel, chromium and iron element. The optical microstructures of Inconel 600 and 690 were shown in Fig. 1a and b, which had been revealed by electrolytic etching in 10% oxalate solution at 5 V for 25 s. The average grain size of Inconel 600 calculated by 200 grains was about 17.4 µm with bulk hardness of 240 HV. Correspondingly, they are 28 µm and 200 HV for Inconel 690. Fretting wear tests were performed on an Optimol SRV IV oscillating friction tester using a ball-on-flat contact configuration at room temperature (25–30 °C) in air condition. The upper specimen was 304SS ball with a diameter of 10 mm, and the lower specimen was cut to a dimension of 15 mm × 15 mm × 2.1 mm then conglutinated to a stainless base. Details of this device had been described in the previous study [11]. The test parameters were set as normal load at 60 N, 82
Wear 416–417 (2018) 81–88
J. Li et al.
Fig. 3. The SEM morphologies as well as oxygen distribution of worn scar surface: (a) Inconel 600 and (b) Inconel 690.
deformed layer (GDL, between dash dot- long dash dot lines). The grain boundaries are the criteria for differentiating TTS and GDL as they cannot be observed in TTS but are deformed and parallel to the fretting direction in GDL. However, it should be mentioned that the WDL thickness of Inconel 600 is about 5 µm, which is larger than that of Inconel 690 (~ 3 µm). But the TTS thickness of Inconel 600 (~ 10 µm) is much thinner than that of Inconel 690 (~ 20 µm). Finally, there is no noticeable difference in GDL thickness between Inconel 600 and 690. Fig. 5 shows the load-indentation depth curve of different microstructure in worn subsurface. The locations tested were about 6 µm (TTS), 26 µm (GDL) and 46 µm (Matrix) from the surface, respectively. The hardness of TTS, GDL and Matrix is 5.33, 4.54 and 3.89 GPa in worn subsurface of Inconel 690. Correspondingly, they are 4.95, and 4.43 and 4.09 GPa in the case of Inconel 600. It is obvious that the hardness decreases from TTS to GDL to Matrix in both cases. In addition, the hardness of TTS and GDL in worn subsurface of Inconel 690 is slightly larger than that of Inconel 600. In order to find out the reason why the thickness and hardness of stratified structures varies with Inconel 690 and 600, the microstructures of GDL and TTS in nanoscale are characterized. Firstly, two TEM samples are obtained from Inconel 690 worn subsurface using FIB. As shown in Fig. 6, sample 1 is taken from the interface of GDL/Matrix, using an original grain boundary (GB) as the reference. Therefore, the section below the original GB is Matrix, the upper is GDL. And sample 2 is taken from the interface of TTS/GDL, using a deformed GB as the reference. Correspondingly, the section below the deformed GB is GDL, the upper section is TTS. To facilitate the following analysis, the directions from Matrix to GDL as well as GDL to TTS are expressed using a to b and b to c, respectively. Fig. 7 shows the characterizations of sample 1 located at the interface of GDL/Matrix. From the bright field TEM (BFTEM) image shown in Fig. 7a, the material above the dash line is platinum, which is deposited and used as the protective layer during the preparation of TEM sample. As a result, the sample could be divided into two parts according to the GB: Grain 1 and Grain 2. Combined with the selected
microscope (TEM) using FEI Tecnai F20. 3. Results Fig. 2a shows the coefficient of friction (COF) as a function of cycle number during fretting wear in the case of Inconel 600 and 690. It could be found that the COF of both cases increases rapidly at the initial stage, and then gradually tends to a steady state with fluctuations. However, although the average COF of both is about 1, the COF of Inconel 600 presents a larger fluctuation than that of Inconel 690. Fig. 2b shows the typical cross-sectional profiles of worn scar along the fretting direction. It could be seen that a U-shape worn scar is produced in both cases. By comparison, the worn scar of Inconel 600 is deeper and wider than that of Inconel 690. Correspondingly, the wear volume is 58.5 × 105 μm3 for Inconel 600 and 47.5 × 105 μm3 for Inconel 690. In fact, the result of parallel experiments indicates that the average wear volume of Inconel 600 is about 1.15 times larger than that of Inconel 690. Fig. 3a and b show the worn scar surface SEM images and the oxygen distribution (the upper right corner) of Inconel 600 and 690, respectively. According to the characterization shown in Fig. 3a and b, in both cases, the oxygen content inside the worn scar is larger than that outside the worn scar, which indicates that the severe oxidation behaviors have occurred on the whole worn scar surface. Then, the enlarged SEM morphologies of position A and B show that both the worn scar surfaces are covered by continuous wear debris. As a result, there is no obvious difference on worn scar surface between Inconel 690 and 600. For further research, the etched worn subsurface SEM images of Inconel 600 and 690 are characterized and shown in Fig. 4a and b, respectively. The locations of worn subsurface shown in Fig. 4a and b correspond to the marks A and B in Fig. 3. It could be seen that there is a stratified structures in the worn subsurface. From the outmost surface to material inside, the structures are wear debris layer (WDL, between dash-dash lines), TTS (between dash-dash dot lines) and general 83
Wear 416–417 (2018) 81–88
J. Li et al.
Fig. 4. The SEM morphologies worn scar subsurface: (a) Inconel 600 and (b) Inconel 690.
brightness of the sample is extremely uneven. According to the block's shape, the sample can be divided into two parts. The left part of dash dot line is equiaxed blocks and the right part is striplike blocks. From the BFTEM image of C position shown in Fig. 8b, it could be found that the striplike blocks correspond to some fine twins and subgrains. They are dozens of nanometers in width and hundreds of nanometers in length. Meanwhile, some dislocation arrays and walls are observed inside the twins and subgrains. As shown in Fig. 8c, the high resolution transmission electron microscope (HRTEM) image of C1 indicates the observed twins could be the mechanical twins due to the dislocations and steps located at the twin boundary, which are the unique characteristics of mechanical twins compared with annealing or growth twins [15]. According to the BFTEM and concentric SADP images of D position shown in Fig. 8d, the equiaxed blocks correspond to randomly oriented equiaxed nano-sized grains. The statistic distribution of equiaxed grain sizes in Fig. 8e shows that the grain size ranges from 20 nm to 90 nm with the average of 47.3 nm.
Fig. 5. The load-indentation depth curve of different microstructure in worn subsurface of Inconel 690 and 600.
4. Discussion In order to increase the thermal conductivity, the thickness of SG tube was reduced from 2 mm to 1.09 mm. As a result, the wear depth was also an importance factor for reliability of SG tube during fretting wear test. As shown in Fig. 2b, both the deeper worn scar and the larger wear volume indicated that Inconel 690 exhibited more excellent wear resistance than Inconel 600. By comparing the morphologies of wear scar surface and subsurface, there was no significant difference on worn scar surfaces as they were both covered by the continuous wear debris as shown in Fig. 3. From the morphologies shown in Fig. 4, fretting wear caused the formation of stratified structures in the worn subsurface, which were WDL, TTS, GDL and matrix from outside into inside, respectively. And it was obvious that Inconel 690 possessed a thinner WDL but a thicker TTS compared with Inconel 600. The thinner WDL of Inconel 690 could be attributed to the lower wear volume for its excellent wear resistance. According to other's previous studies, the thicker dense and continuous WDL could be more effective to separate the contacting surface and act as a lubricant. To some extent, it has benefits for the improvement of material's wear resistance [16–18]. However, compared with Inconel 600, Inconel 690 with thinner WDL but excellent wear resistance indicated that the wear resistance was not primarily dependent on WDL in this research. Therefore, the reason why Inconel 690 presented more excellent fretting wear resistance than Inconel 600 could depends mainly on TTS as it was directly related to material loss, which could be explained as follows. On one hand, it had been reported that the wear resistance was proportional to the hardness of material within a certain range [19]. The TTS's hardness of Inconel 690 was greater than that of Inconel 600 as shown in Fig. 5, resulting in an increase in wear resistance. On the other hand, the increasing in the thickness of the deformation layer would reduce the transfer of stress
Fig. 6. The schematic view of FIB sampling position at Inconel 690 worn subsurface.
area diffraction pattern (SADP) of A and B as shown in Figs. 7b and 7c, the high-density tangled dislocations are formed in Grain 1. Meanwhile, the multiple twins associated with high-density tangled dislocations are formed in the Grain 2. The length of twins is up to several micrometers, but the width of twins is only hundreds of nanometers. From the dark field TEM (DFTEM) of A and B shown in Figs. 7d and 7e, the brightness in Grain 1 or Grain 2 are inconsistent, which indicates that the grains have been subdivided into several regions with different crystallographic orientations. Fig. 8 shows the characterizations of sample 2 located at the interface of TTS/GDL. From the DFTEM image shown in Fig. 8a, the
84
Wear 416–417 (2018) 81–88
J. Li et al.
Fig. 7. The characterization of sample 1 located at the interface of GDL/Matrix: (a) overall BFTEM image, (b) SADP of A position, (c) SADP of B position, (d) DFTEM of A position and (e) DFTEM of B position.
85
Wear 416–417 (2018) 81–88
J. Li et al.
Fig. 8. The characterizations of sample 2 located at the interface of TTS/GDL: (a) the overall DFTEM image, (b) the BFTEM of C position, (c) the HRTEM of C1 position, (d) the BFTEM as well as SADP images of D position and (e) the statistics of grain size in TTS.
gradually from contact surface to material inside. In order to accommodate plastic strains in polycrystalline materials, various dislocation activities were normally motivated in low strain conditions [21,22]. As a result, there were plenty of tangled dislocations in matrix (I stage, Fig. 7-Grain 1). Secondly, twinning was an important way of plastic deformation, which often acted as a supplement when dislocation slips was not easy to occur [23,24]. In addition, experimental results indicate that the critical stress for twinning was larger than dislocation slip [25]. Thence, many mechanical twins associated with dislocations were formed in coarse grain (II stage, Fig. 7-Grain 2). Furthermore, the twins would continue to undergo twin deformation with the further increase of strain, resulting in a reduction of the thickness and length of twins
from the surface to the substrate [20]. This would also enhance the material's wear resistance. As mentioned above, the wear resistance was closely related to the thickness and hardness of TTS. As a result, the relationship of TTS's thickness and hardness with nickel-based alloy would be discussed based on the formation mechanism of TTS. The first was about the formation mechanism of TTS. According to the microstructure characterization shown in Figs. 7 and 8, the transformation from matrix to TTS could be proposed using the schematic diagram shown in Fig. 9. According to previous research results of others, the formation mechanism of TTS could be analyzed as follows. Firstly, it is well known that the strain could be generated during fretting wear and decrease 86
Wear 416–417 (2018) 81–88
J. Li et al.
Fig. 9. Schematic diagram for TTS formation.
COF on worn scar surface of Inconel 600 was greater than that of Inconel 690.
and subgrains (III stage, Fig. 8b). In addition, including the strain, friction heat was also generated on contact surface during fretting wear. According to the recrystallization theory, the increase of temperature could cause the formation of dislocation walls as shown in Fig. 8b. With the further increase of strain and temperature, the tangled dislocation and dislocation walls could act as nuclei. Finally, the striplike twins and subgrains transformed into equiaxed grain (IV stage, Fig. 8d) as friction heat and strain increased and reached to the threshold value for dynamic recrystallization (DC). In summary, in nanoscale, GDL was mainly composed of striplike twins and subgrains, while TTS consisted of equiaxed nanosized grains. The transformation mechanism from the matrix (coarse grains) to TTS (nano-sized grains) was DC. The formation of TTS was in the form of DC according to the above discussion. Then, the relationship between DC and nickel-based alloy should be revealed. It is well known that the DC is closely related to material's stacking fault energy (SFE). Materials with lower SFE were usually prone to DC as the dislocation climb and cross slip were difficult to occur [26–28]. In addition, the research about the effect of chromium content on SFE showed that SFE rapidly decreased with the increase of chromium content in nickel-base alloys [29]. From the chemical composition shown in Table 1, the chromium content of Inconel 690 was about 30 wt%, which was as twice as that of Inconel 600. In short, the higher chromium content led to the lower SFE, which prompted the occurrence of DC and eventually created a thicker and harder TTS in Inconel 690 worn subsurface. As a result, compared with Inconel 600, Inconel 690 with thicker TTS exhibited more excellent wear resistance. Furthermore, it should be mentioned that the reduction of grain size was conducive to the formation of TTS due to the increasing of bulk hardness, which had been proposed in our previous work [11]. However, the bulk hardness of Inconel 690 was smaller than that of Inconel 600, but had more excellent wear resistance in present study. As a result, compared with grain size, the Cr content could present a greater impact on the thickness of TTS. As for the COF shown in Fig. 2a, it was reported that the fluctuation of COF was closely related to the formation of wear debris and their ejection from the contact area [30]. Compared with Inconel 690, more wear debris generated during fretting wear of Inconel 600, which could be deduced from the larger wear volume. Therefore, the fluctuation of
5. Conclusion The fretting wear behaviors of Inconel 600 and 690 in the atmosphere were investigated using the 3D surface profiler, SEM, EDX and TEM. The observation and analysis led to the following conclusions: (1) Fretting wear caused the formation of stratified structures in the worn subsurface of Inconel 690 and 600, which were WDL, TTS, GDL and matrix from outside into inside, respectively. (2) Compared with Inconel 600, a thinner WDL associated with a thicker and harder TTS were formed in worn subsurface of Inconel 690. (3) The transformation mechanism from the matrix (coarse grains) to TTS (nano-sized grains) was dynamic recrystallization. (4) Inconel 690 exhibited more excellent wear resistance than Inconel 600, which was mainly attributed to the thicker and harder TTS in the worn subsurface of Inconel 690 for its high chromium content. Acknowledgements This work was supported by the National Key Basic Research Program of China (973 Program) (No. 2011CB610504) and the Guangdong Province Science and Technology Plan (No. 2017B090903005). References [1] M.H. Attia, Fretting fatigue and wear damage of structural components in nuclear power stations-fitness for service and life management perspective, Tribol. Int. 39 (2006) 1294–1304. [2] F.M. Guérout, N.J. Fisher, Steam generator fretting-wear damage: a summary of recent findings, J. Press. Vessel. Technol. 121 (1999) 304–310. [3] M.K. Lim, S.D. Oh, Y.Z. Lee, Friction and wear of Inconel 690TT and Inconel 600 for steam generator tube in room temperature water, Nucl. Eng. Des. 226 (2003) 97–105. [4] Y. Yoon, I. Etsion, F.E. Talke, The evolution of fretting wear in a micro-spherical contact, Wear 270 (2011) 567–575.
87
Wear 416–417 (2018) 81–88
J. Li et al.
steel contacts, Tribol. Int. 44 (2011) 1452–1460. [19] Y.S. Zhang, Z. Han, K. Lu, Fretting wear behavior of nanocrystalline surface layer of copper under dry condition, Wear 265 (2008) 396–401. [20] S. Fayeulle, P. Blanchard, L. Vincent, Fretting behavior of titanium alloys, Tribol. T. 36 (1993) 267–275. [21] B.J. Griffiths, D.C. Furze, Tribological Advantages of White Layers Produced by Machining, J. Tribol. 109 (1987) 338–342. [22] W.L. Chan, M.W. Fu, J. Lu, The size effect on micro deformation behavior in microscale plastic deformation, Mater. Des. 32 (2011) 198–206. [23] Z. Wu, H. Bei, F. Otto, G.M. Pharr, E.P. George, Recovery, recrystallization, grain growth and phase stability of a family of FCC-structured multi-component equiatomic solid solution alloys, Intermetallics 46 (2014) 131–140. [24] C. Zheng, D. Raabe, Interaction between recrystallization and phase transformation during intercritical annealing in a cold-rolled dual-phase steel: a cellular automaton model, Acta Mater. 61 (2013) 5504–5517. [25] Y.T. Zhu, X.Z. Liao, X.L. Wu, J. Narayan, Grain size effect on deformation twinning and detwinning, J. Mater. Sci. 48 (2013) 4467–4475. [26] S.L. Guo, D.F. Li, Z.G. Wu, Constitutive relationship and critical condition for dynamic recrystallization of inconel 625 during hot deformation, Adv. Mater. Res. 538–541 (2012) 1240–1244. [27] M.S. Chen, Y.C. Lin, X.S. Ma, The kinetics of dynamic recrystallization of 42CrMo steel, Mater. Sci. Eng. A 556 (2012) 260–266. [28] N.R. Tao, X.L. Wu, M.L. Sui, Grain refinement at the nanoscale via mechanical twinning and dislocation interaction in a nickel-based alloy, J. Mater. Res. 19 (2004) 1623–1629. [29] D.M. Symons, Hydrogen embrittlement of Ni-Cr-Fe alloys, Metall. Mater. Trans. A 28 (1997) 655–663. [30] N. Diomidis, S. Mischler, Third body effects on friction and wear during fretting of steel contacts, Tribol. Int. 44 (2011) 1452–1460.
[5] J.Y. Yun, M.C. Park, G.S. Shin, Effects of amplitude and frequency on the wear mode change of Inconel 690TT SG tube mated with SUS 409, Wear 313 (2014) 83–88. [6] Z.H. Wang, Y.H. Lu, J. Li, T. Shoji, Effect of pH value on the fretting wear behavior of Inconel 690 alloy, Tribol. Int. 95 (2016) 162–169. [7] J. Li, Y.H. Lu, Effects of displacement amplitude on fretting wear behaviors and mechanism of Inconel 600 alloy, Wear 304 (2013) 223–230. [8] X.Y. Zhang, P.D. Ren, J.F. Peng, Fretting wear behavior of Inconel 690 in hydrazine environments, Trans. Nonferrous Met. Soc. China 24 (2014) 360–367. [9] Y. Lee, H. Kim, A comparative study on the fretting wear of steam generator tubes in Korean power plants, Wear 255 (2003) 1198–1208. [10] L. Xin, B.B. Yang, J. Li, Z.H. Wang, Y.H. Lu, T. Shoji, Microstructural evolution of subsurface on Inconel 690TT alloy subjected to fretting wear at elevated temperature, Mater. Des. 104 (2016) 152–161. [11] J. Li, Y.H. Lu, H.Y. Zhang, L. Xin, Effect of grain size and hardness on fretting wear behavior of Inconel 600 alloys, Tribol. Int. 81 (2015) 215–222. [12] M.A.L. Tobi, J. Ding, G. Bandak, S.B. Leen, P.H. Shipway, A study on the interaction between fretting wear and cyclic plasticity for Ti-6Al-4V, Wear 267 (2009) 270–282. [13] B. Yao, Z. Han, K. Lu, Correlation between wear resistance and subsurface recrystallization structure in copper, Wear 3 (2012) 438–445. [14] J. Li, B.B. Yang, Y.H. Lu, L. Xin, Z.H. Wang, T. Shoji, The effects of electrochemical polarization condition and applied potential on tribocorrosion behaviors of Inconel 690 alloys in water environment, Mater. Des. 119 (2017) 93–103. [15] L. Lu, Y. Shen, X. Chen, L. Qian, K. Lu, Ultrahigh strength and high electrical conductivity in copper, Science 304 (2004) 422–426. [16] M. Verenberg, G. Halperin, I. Etsion, Different aspects of the role of wear debris in frettingwear, Wear 252 (2002) 902–910. [17] J. Warburton, The fretting of mild steel in air, Wear 131 (1989) 365–386. [18] N. Diomidis, S. Mischler, Third body effects on friction and wear during fretting of
88